bio-polyamides based on renewable raw materials · 2017. 8. 29. · doi 10.1007/s10973-015-4929-x....
TRANSCRIPT
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Bio-polyamides based on renewable raw materials
Glass transition and crystallinity studies
Joanna Pagacz1,2 • Konstantinos N. Raftopoulos1,3 • Agnieszka Leszczyńska1 •
Krzysztof Pielichowski1
Received: 3 March 2015 / Accepted: 14 July 2015 / Published online: 5 September 2015
� The Author(s) 2015. This article is published with open access at Springerlink.com
Abstract Structural characterization of a series of novel
bio-polyamides based on renewable raw materials—PA
4.10, PA 6.10, PA 10.10, and PA 10.12—was performed by
Fourier transform infrared spectroscopy (FTIR) and wide-
angle X-ray diffraction (WAXD). Infrared spectra and the
WAXD patterns indicate the coexistence of different
crystalline forms, a- and c-triclinic and b-pseudohexago-nal. Thermal properties in the glass transition (Tg) and
melting region were then investigated using temperature-
modulated DSC (TOPEM� DSC). The melting point (Tm)
was found to increase with increasing amide/methylene
ratio in the polymer backbone, which is consistent with the
increase in linear density of hydrogen bonds. Studies on the
molecular dynamics by dynamic mechanical analysis show
three distinct regions associated with the c- and theb-relaxation and the dynamic glass transition. TOPEM�
DSC data reveal that at low frequency/long timescales, the
materials with significantly different amide/methylene
ratios have similar segmental dynamics.
Keywords Biopolymers � Glass transition � Moleculardynamics � Differential scanning calorimetry � WAXS �Polymorphism
Introduction
Biopolymers produced from renewable raw materials pro-
vide an environmentally friendly alternative to conven-
tional petroleum-based polymers and exhibit new
interesting properties, such as low water uptake, high
mechanical resistance, high melting point, and crystal-
lization rate [1–4]. Vegetable oils derived from inedible or
toxic plants are good alternative chemical feedstock due to
their reactivity (numerous active chemical sites) and bio-
compatibility. They are vastly used for polyurethanes and
epoxies; however, durable biothermoplastics, such as bio-
polyamides (bio-PAs), based on renewable sources have
high growth potential as they show promising mechanical
and thermal properties [5–7].
Bio-polyamides are synthesized from two or more
monomers or comonomers, which belong to amino acid,
cyclic amide (lactam), dicarboxylic acid, and diamine
families. Fatty acids present in vegetable oils can be con-
verted by simple reactions into suitable bifunctional
monomers for the production of polyamides by a simple
polycondensation process. Among most known biomono-
mers are 11-aminoundecanoic acid and sebacic acid pro-
duced by conversion of ricinoleic acid derived from castor
oil—Scheme 1.
Typically, raw castor oil is hydrolyzed to give ricinoleic
acid, which is then converted to sebacic acid in a reaction
with potassium or sodium hydroxide at high temperature
[8]. The 1,12-dodecanedioic is most often prepared from
petrochemical butadiene, but potentially it can be obtained
by an x-oxidation process of lauric acid catalyzed byyeast strain [9]. Undecane-1,11-dicarboxylic acid and
11-aminoundecanoic acid were reported as potential
monomers, too [10]. The diamines for the production of
bio-polyamides, e.g., 1,5-diaminopentane (cadaverine),
& Joanna [email protected]
1 Department of Chemistry and Technology of Polymers,
Cracow University of Technology, ul. Warszawska 24,
31-155 Kraków, Poland
2 Present Address: Wroclaw Research Centre EIT?, ul.
Stabłowicka 147, 54-066 Wrocław, Poland
3 Present Address: Physik-Department, Fachgebiet Physik
weicher Materie, Technische Universität München, James-
Franck-Str. 1, 85748 Garching, Germany
123
J Therm Anal Calorim (2016) 123:1225–1237
DOI 10.1007/s10973-015-4929-x
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pentamethylenediamine, or tetramethylenediamine, are
naturally occurring substances or they can be produced by
microbial biosynthesis (e.g., by decarboxylation of amino
acids (lysine, ornithine) and polymerization with sub-
stances from microbial fermentation (such as succinate)
[11–13].
The other way to produce polyamides (PAs) is based on
bifunctional monomers; for example, Yamano et al. [14]
reported on the preparation of branched polyamide 4 from
c-aminobutyric acid originated from biomass fermentation.The final product was, unlike other polyamides,
biodegradable by bacteria isolated from activated sludge.
Other bifunctional monomers can be synthesized by com-
plementary hydrobromination of fatty acid methyl esters
and further modification [15]. The polyamides with dif-
ferent chain lengths between amide groups (ratio of
methylene and amide groups) show different hydrogen
bond densities and thus different water absorption. In
general, the longer carbon chain results in higher impact
strength, higher solvent resistance, and lower water
absorption [16, 17]. These characteristics bring possibilities
to design novel biobased materials with specially tailored
properties.
The degree of crystallinity of polymers, including bio-
polyamides, determines their physical and mechanical
properties. For semicrystalline materials, a two-phase model
of the structure was proposed with a crystalline and an
amorphous part, but also a third phase can exist. This phase,
the so-called rigid amorphous fraction (RAF), is located at
the interfaces between the crystalline and amorphous
regions, and is characterized by intermediate properties
between two main phases and reduced mobility [18].
Hydrogen bonding plays a significant role in the struc-
ture of polyamides, especially in the 2N 2(N ? 1) poly-
amides, i.e., polyamides with equal amounts of methylene
groups in the diamine and diacid segment. There are two
types of hydrogen bonds that may occur in polyamides—
the amide–amide bond and the amide–ester hydrogen bond.
Crystals of polyamides were found to be chain-folded and
arranged in sheets held together by hydrogen bonds [18].
There are two possible hydrogen-bonded schemes for
chain-folded sheets of PAs—the alternating and the pro-
gressive one, and they can stack with progressive shearing,
yielding the so-called a-phase or with alternating shearing,which is known as the b-phase. In polyamides2N 2(N ? 1), both types of sheet stacking were found,
which gives four possible crystal structures [19]. The XRD
pattern for these polyamides shows characteristic diffrac-
tion peaks related to the interchain and intersheet distance.
Moreover, polymorphism and solid transitions can be
observed in polyamides. The former effect is related to the
formation of different crystal structures, i.e., the planar a-and the pleated c-structures, while the latter is an effect ofcrystal phase transformation prior to melting [20].
The purpose of this work was to study the crystallinity
and molecular dynamics of four novel biobased poly-
amides with variable ratio of methylene and amide groups
in the polymer backbone. A novel thermal technique of
stochastic modulated-temperature differential scanning
calorimetry (DSC TOPEM�) is employed for the study of
both phenomena, in comparison and complementarity with
well-established techniques of wide-angle X-ray diffraction
(WAXD) and dynamic mechanical analysis (DMA). In
addition, Fourier transform infrared spectroscopy (FTIR)
was used as a complementary method for the study of
crystalline and amorphous fraction in polyamides.
Experimental
Materials
The specimens in this study are partly or fully bioresourced
polyamides, produced by DSM (PA 4.10, EcoPaXX�
Q170MS) and Evonik (Vestamid Terra� PA 6.10, PA
10.10, PA 10.12) by polycondensation of adequate diamine
and diacid (Scheme 1), with the specifications presented in
the Table 1 [16, 21].
The biobased carbon content (BCC) was determined by
the producer according to the standard ASTM 6866 and is
well above 50 mass% [22]. The PA 10.12 differs in the
BCC values, because one of the monomers used for its
synthesis, 1,12-dodecanedioc acid, can be sourced both by
petroleum production route and via bioproduction from
palm kernel oil.
Methods
Processing
For the FTIR, DMA, and WAXD studies, injection-molded
samples were used. The pellets of bio-polyamides after
PALM KERNEL OIL
dodecanedioic acid
PA 10.12
PA 10.10
PA 4.10 PA 6.10
+HMDA +TMDA
decamethylene diamine
CASTOR OIL
ricinoleic acid
sebacic acid
Scheme 1 Schematic route of bio-polyamides production fromnatural oils
1226 J. Pagacz et al.
123
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drying (70 �C, 8 h) were processed on an injection moldingmachine ZAMAK WT 12 at temperature 30 �C above theirrespective melting points. The mold temperature was kept
at 90 �C below the respective crystallization temperature.
Structural characterization
Fourier transform infrared spectroscopy (FTIR) Fourier
transform infrared (FTIR) spectra were recorded with a
PerkinElmer Spectrum 65 spectrophotometer in ATR mode
with C/ZnSe crystal; eight scans with resolution 4 cm-1 in
the wave number range of 600–4000 cm-1 were
performed.
Wide-angle X-ray diffraction (WAXD) A Bruker D2Phaser
diffractometer with CuKa-radiation (k = 0.154 nm) wasused to study the crystalline structure of the bio-polyamides.
The X-ray diffraction measurements were taken in the
reflection mode at room temperature on 1-mm-thick samples
(solid bars) produced by injection molding. Bragg’s equa-
tion, d = nk/2sinh, was applied to calculate the interlayerspacing (d001) of the samples by using 2h-values from theWAXD patterns. Deconvolution of the peaks observed in
WAXD analysis was performed with the software Grafity
version 0.4.5. The analysis yielded for each reflection peak
its position in the 2h domain, its area, and its full width at halfmaximum (FWHM).
Thermomechanical characterization
Modulated-temperature DSC (TOPEM�) The DSC with
stochastic temperature modulation, provided by Mettler
Toledo and branded as TOPEM�, analyzes the dynamic
behavior of materials during one single measurement with
multifrequency modulation [23]. Step-like temperature
perturbations with fixed amplitude and random duration
are superimposed to the conventional linear heating, and
the heat flow response is recorded. Upon a Fourier
transformation, it is possible to calculate the quasi-static
heat capacity, cp,0, and the frequency-dependent complex
heat capacity, c�p (x), at different frequencies and, thus,
provide information on the dynamic response of the
thermal events. In a more conventional approach of
modulated DSC, the stochastic heat flow response may be
deconvoluted to its reversing and the non-reversing
components, which are related to the thermodynamic and
kinetic phenomena, respectively. The method needs to run
at sufficiently low heating rate (0.01–2 �C min-1) andsmall values of temperature perturbation dT0 (1 mK–0.5 K).
In this work, the melting and the glass transition regions in
bio-polyamides were studied in two separate stochastic
modulated runs, on fresh as-received samples, each upon a
preparatory protocol to remove thermal history and provide
good contact with the crucible. Thus, the material was heated
up to 280 �C, maintained for 3 min, and then cooled to-10 �C, in a conventional run at a rate of 20 �C min-1. Thestochastic modulation runs were performed between -10
and 90 �C for the study of the glass transition region and invarying temperature ranges between 200 and 290 �C, basedon preliminary measurements on each sample, for the study
of crystallinity. The underlying heating rate was in both
cases 0.5 �C min-1, the amplitude of thermal pulses was setat 1 K, and each pulse had a random duration in the range of
15–30 s (switching time range). All measurements were
taken in argon. Typical sample mass was about 20 mg, and
20-ll aluminum crucibles were used.
DMA analysis Dynamic mechanical analysis (DMA
242C Netzsch) was carried out on the bars obtained by the
injection molding process (area 10 9 26 mm2, thickness
1 mm) in three-point bending mode (20-mm sample
holder) under maximal dynamic force of 6 N, constant part
of static force of 0.3 N, and static force 10 % higher than
dynamic. The maximal amplitude was set to 50 lm. Atemperature scan from -160 to 150 �C was performedat 2 �C min-1 and with perturbation frequencies of 1,2.5, 5, and 10 Hz. Liquid nitrogen was used as cooling
medium and dry air as purge. The viscoelastic parameters
of storage modulus (E0), loss modulus (E00), and mechan-ical loss tangent (tand) were recorded as a function oftemperature.
Table 1 Characteristics of bio-polyamides
Bio-PA Structure CONH/CH2 VN/cm3 g-1 BCC/%
PA 4.10 [–NH(CH2)4NHCO(CH2)8CO–] 0.17 170 70
PA 6.10 [–NH(CH2)6NHCO(CH2)8CO–] 0.14 220 63
PA 10.10 [–NH(CH2)10NHCO(CH2)8CO–] 0.11 220 100
PA 10.12 [–NH(CH2)10NHCO(CH2)10CO–] 0.10 220 45–100
VN Viscosity number (ISO 307), BCC biobased carbon content (ASTM 6866)
Renewable-sourced bio-polyamides 1227
123
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Results and discussion
Structural characterization (FTIR)
Figure 1 shows the FTIR spectra of the bio-polyamides.
The bands registered for bio-PAs with their respective
assignments are summarized in Table 2.
The infrared spectra show characteristic bands for aliphatic
polyamides, corresponding to methylene (C–H stretching
2920/2851 cm-1), amide I (C–O stretch 1637 cm-1), amide
II (in-plane N–H deformation with CO and CN stretches,
ca. 1535 cm-1), and ester group vibrations (1741 cm-1),
respectively. The N–H stretching mode appears as a single
band at 3300 cm-1, which indicates strong hydrogen bond-
ing, and it is rather broad (Fig. 1), which may suggest that
more types of hydrogen bonds with different bond lengths
were formed [24]. Moreover, the shape (no shoulders) of the
amide I (1637 cm-1) and II bands suggests also that some of
the carbonyl groups are hydrogen bonded [25]. This is in
agreement with the literature data, indicating that in poly-
amides two types of hydrogen bonds can be observed,
amide–amide and ester–amide, depending on the position
of the ester and the amide group along the polymer
chain [26].
The bands in the region of 1200–900 cm-1 are related to
the different crystalline structures in the polyamides. The
a-structure should give the absorbance at 1200 and930 cm-1, and the c-crystalline phase at 973 cm-1 [27].On the spectra shown in Fig. 1, the bands around 1190 and
942–948 cm-1 (C–CO stretching ‘‘crystalline’’ band)
confirm the predominant existence of the a-crystallinestructure in all tested polyamides. In general, clear differ-
ences can be observed in the region 1148–1200 cm-1.
All polyamides show the ‘‘crystalline’’ band at ca.
1190 cm-1; however, its intensity for PA 6.10 is lower
than for the other materials. Interestingly, a new band at
1179 cm-1 appears for PA 6.10, which is also related to the
crystalline domain, but it was found as non-dependent on
the crystallinity [27]. This band was also observed for PA
10.10 and 10.12, but at slightly lower frequency of
1165 cm-1. Galimberti et al. [29] attribute the broad band
at 1123 cm-1 to the amorphous phase related to the chain
defects, and it was observed for all bio-PAs.
Moreover, the weak band at 877 cm-1 in the spectrum
of polyamide 6.10 can be related both to the amorphous
region presence and peroxide C–O–O– vibrations, coming
from the partial oxidation during processing.
Crystalline structure (WAXD)
The WAXD patterns of bio-polyamides are presented in
Fig. 2.
The characteristic reflection (100) located at about
2h = 20� and corresponding to the interchain distance inthe crystalline structure of the a-crystal form was observedas predominant phase for all samples. The 001 and 002
diffraction peaks at 2h * 7 and 10� are related to thelength of the chemical repeat unit, but less visible and
rather broad.
The relative intensity of the 001 and 002 reflections was
found to be governed by the distance between the carbonyl
groups in the two alkane segments—if the distance is rather
equal, the 001 reflection would be weak because of the
destructive interference [34].
A strong reflection located at 23� and assigned to010/110 crystallographic planes was observed for bio-
polyamides, except PA 6.10. Jasińska et al. [24] found
similar results for polyamide 4.10 and its copolymers, and
they described the peak at 2h = 23� as characteristic of thec-crystalline phase. On the contrary, for PA 6.6, Liu andcoworkers assign this reflection as one of the distinctive
features of the a-phase in polyamides and state that poly-amides present usually more stable a-phase rather than cones in XRD [20]. In the c-form phase of the polyamide,hydrogen bonds take position between parallel macro-
chains, while in a-form the antiparallel arrangement ispresent, and the latter is more stable because of shorter
hydrogen bonds [35].
The scattering patterns for PAs 4.10, 10.10, and 10.12
are comparable, but they present differences in the relative
intensity of 100 and 010/110 reflections. This means that
the intersheet distance is not affected by the backbone
length in the polyamides, while the higher concentration of
the amide groups (CONH/CH2 ratio) causes visible chan-
ges in the intensity of the 100 peak in comparison with the
010/110 reflection. In PA 4.10, the CONH/CH2 ratio is
higher than in PA 10.12, and the intensity of the 001 band
increases in relation to 010/110 one. Interestingly, for PA
3500 3000 2500 1800 1600 1400 1200 1000 800
Wavenumber/cm–1 Wavenumber/cm–1
PA 4.10PA 6.10PA 10.10PA 10.12
3300 cm–1
1741 cm–1
1190 cm–1
940 cm–1
877 cm–1
*1179 cm–1
Abs
orba
nce/
a.u.
Abs
orba
nce/
a.u.
*
Fig. 1 FTIR spectra of bio-polyamides in different absorbanceregions
1228 J. Pagacz et al.
123
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10.12, the intensities of the 100 and 010/110 reflections are
approximately the same, and a weak band can be seen at
ca. 21�. This observation may indicate the coexistence ofdifferent crystalline forms, a- and c-triclinic and b-pseu-dohexagonal structure. The b-mesomorphic structure wasidentified, e.g., for PA 6 [36]; however, other studies
[34, 37] show that only one crystalline phase exists at room
temperature in PA 10.12. These discrepancies may come
from the different thermal history of the material, resulting
in different crystallization conditions.
Meanwhile, the WAXD pattern for PA 6.10 (Fig. 3)
differs in the region of 2h of 21�–25�, and this suggests thatonly the a-crystalline form is present. However, it is knownthat the polymorphic structures in polyamides come from
different spatial arrangement in the hydrogen bondings;
thus, this may be also related to the less perfect crystalline
structure of PA 6.10 [35]. On the other hand, FTIR spectra
for PA 6.10 show an extra ‘‘amorphous’’ small peak (at
877 cm-1) and a new ‘‘crystalline’’ one (at 1179 cm-1),
while the other ‘‘crystalline’’ markers were diminished.
This corresponds well with the lack of one missing crys-
talline diffraction peak in XRD profiles and may indicate
that the ‘‘crystalline’’ FTIR band emerging for PA 6.10 at
1179 cm-1 (Table 2) corresponds to other a-crystallinestructure.
It was reported for PA 6.10 and PA 10.10 that they may
exist in two triclinic structures, the a- and the b-structures[38], while for PA 10.12 only one crystalline phase was
found [37]. Here, the WAXD results show that bio-PAs,
except PA 6.10, may form different crystalline structures
Table 2 Assignments of the FTIR bands for bio-polyamides
Reference band Bio-PAs Band assignment
4.10 6.10 10.10 10.12
General
3300 [28] 3300 3300 3303 3300 N–H stretch hydrogen bonded
3070 [28] 3072 3072 3072 3084 N–H stretch and amide II overtone
2935 [28] 2920 2920 2920 2920 Asymmetric methylene stretch
2860 [28] 2851 2851 2851 2851 Symmetric methylene stretch
1741 [25] 1741 1741 1741 1740 C=O stretch (ester)
1660 [28] 1632 1637 1637 1637 Amide I: C=O stretch
1530 [28] 1535 1535
1541
1535 1537
1542
Amide II: N–H in-plane bending
coupled with C–N and C–O stretch
1170 [29] – 1179 1165 1165 CO–NH skeletal, crystalline
1123 [29] 1123 vw 1122 1122 1122 mi Amorphous
a-structure
1466 [30] 1466 1466 1466 1466 CH2 scissoring not adjacent to the
amide group
1416 [31] 1419 mi 1419 1419 mi 1420 CH2 scissoring
1373 [31] 1373 vw 1372 1372 1370 Amide III and CH2 wagging
1262 [28] – 1279 – 1273 Amide III
1200 [27] 1192 1192 vw 1191 1188 CH2 twist-wagging
959 [27] 959 vw 959 vw 959 vw 960 vw CO–NH in-plane (shoulder)
936 [29] 947 936 942 941 Vibrations of the N-vicinal CH2group coupled
amide III ‘‘crystalline band’’
a- and c-structures
730 [28] 720 vw 721 721 720 Rocking mode of CH2
c-structure
1439 [32] 1437 1437 1437 1437 mi CH2 scissors vibration
1369 [31] 1358 1363 1360 1370 CH2 twist-wagging
1329 [29] – – – 1336 w C–H deformation
1236 [33] 1231 1237 1237 1237 CH2 twist-wagging
1255 [30] 1255 – – – Skeletal C–C stretch
976 [27] 972 vw 972 vw – 972 vw CO–NH in-plane
vw Very weak, w weak, m medium, mi more intensive
Renewable-sourced bio-polyamides 1229
123
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depending on the CONH/CH2 ratio and thus the hydrogen
bonding density. In PA 6.10, only one type of crystalline
domain was found; however, the method of specimen
preparation should be taken into account.
In general, crystallization depends on the energies of the
amides and the methylene groups, interchain bonding and
chain folding, which are time–temperature characteristics.
Quenching of the material will favor the formation of less
stable c-phase, whereas a slow cooling procedure will leadmainly to the formation of more stable crystalline fraction
[39]. Considering that the same supercooling state was
provided during the processing and the fact that still
crystallization may occur at different cooling rates, it can
be concluded that the processing of bio-polyamides results
in the formation of different crystalline forms.
In order to quantify these observations and calculate the
apparent degree of crystallinity (vc,WAXD), we followed afitting approach. An adequate fit was achieved with a sum
of four Lorentzians to account for crystalline reflections
and three Lorentzians more to account for the amorphous
halo and the smectic phase. The peak positions of the fit-
ting procedure and the spacings as calculated by Bragg’s
law are reported in Table 3. Within the experimental
inaccuracy, the spacings do not show any significant dif-
ferences according to the literature for PAs [18, 19, 24, 40].
The amorphous regions are reflected on the smectic peak at
2h between 18� and 20� and the amorphous halo at ca. 39�.Moreover, the reflections at *21� can be assigned as c-phase reflections, formed during quenching from the melt
state [35, 41]. The smectic phase found around 20� wasalso reported for PA 11 using the temperature-dependent
X-ray measurements [42, 43]. After deconvolution, a single
reflection related to the a-phase was found for PA 6.10(Fig. 3) at 2h = 20.9� and another reflection at 2h = 22�,representing the mesomorphic b-form, and this stays ingood agreement with the reference data [34, 41].
Concluding, we found that for polyamides used in this
study, three crystal phases coexist at room temperature,
except PA 6.10, where only the more perfect a-phase andthe pseudohexagonal b-phase are present.
Moreover, we calculated the apparent degree of crys-
tallinity, vc,WAXD, as the ratio of the areas enclosed by thecrystalline peaks (RAcrystal) to the sum of the areas of thecrystalline peaks and the amorphous profile (RAamorphous) [44]:
vc;WAXD ¼RAcrystal
RAcrystal þ RAamorphous
� �� 100 ð1Þ
There is no clear dependence of the crystallinity in rela-
tion to the amide/methylene ratio; however, the highest
crystallinity was found for PA 10.12, in which well-defined
structure was observed. The values will be discussed later in
relation to the degree of crystallinity evaluated from
TOPEM� DSC measurements. However, it was found for
polyamides 2N (2 N ? 1) that as a length of alkane seg-
ments increases, the change in the sheet structure occurs,
resulting in change in the chemical nature of lamellar sur-
face [19]. Thus, the crystallinity degree will be affected by
the consecutive alkane segments and the CONH/CH2 ratio.
5 10 15 20 25 30
5 10 15 20 25 30
2θ/°
Inte
nsity
/a.u
.In
tens
ity/a
.u.
PA 4.10PA 6.10PA 10.10PA 10.12
PA 4.10 experimental100 (α-structure)Undefined001002Smectic010/110AmorphousFit
100
010/110
001 002
2θ/°
Fig. 2 Up WAXD patterns of all bio-polyamides under investigation.Down A representative profile fitting result for the PA 4.10
8000
7000
6000
5000
4000
3000
2000
1000
0
10 20 30 40 50 60
Inte
nsity
/a.u
.
PA 6.10 experimental002Smectic
AmorphousAmorphousFit
100β-phase
2θ/°
Fig. 3 Fitting profile of WAXD pattern for PA 6.10
1230 J. Pagacz et al.
123
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Melting and crystallization (TOPEM� DSC)
An examination of four bio-polyamides for the melting
behavior and the crystallinity was conducted by stochastic
modulated DSC (TOPEM�). We used the formalism of the
reversing and non-reversing heat flow. The corresponding
curves are shown in Fig. 4.
The enthalpy of fusion, DHf, was calculated as the dif-ference in the endothermic melting from the reversing heat
flow component (DHrev) and the exothermic signal on thenon-reversing heat flow profile (DHnon-rev): DHf =DHrev - DHnon-rev. The obtained value was further used tocalculate the crystallinity degree according to:
vc;DSC ¼DHfDH0f
� 100 ð2Þ
where DHf is the experimental enthalpy of fusion (J g-1)
and DH0f is the reference value of enthalpy of fusion forpurely crystalline polyamides. The reference values are 254
and 244 (J g-1) for polyamides 6.10 and 10.10, respec-
tively [38, 45]. To the best of our knowledge, there is no
available data for other polyamides. However, considering
the chemical similarity of the tested materials, the DH0fvalues for PAs 6.10 and 10.10 were taken as reference
values for their closest counterparts, i.e., for PA 4.10 and
PA 10.12, respectively. The melting and crystallization
points, the associated enthalpies for bio-polyamides, are
summarized in Table 4.
The high melting point for PAs was found as the result
of the presence of hydrogen bonds which influence the
macrochain mobility below Tm and the order of macro-
chains in the molten state. It was reported that hydrogen
bonds may order the crystalline nuclei in the melt and thus
influence the subsequent crystallization [39].
The reversing signal contains the melting of outer layers
of crystallites. Interestingly, the melting peaks are typically
complex with more than one component, corresponding to
the two different types of crystal forms as observed by
WAXD. In agreement with this interpretation, the revers-
ing heat flow curve of the PA 6.10 shows only one peak
corresponding to the a-phase.The multiple peak of melting in polyamides can be
assigned to the fusion of imperfect crystals with lower thermal
stability, secondary crystallization upon heating, and remelt-
ing of more perfect crystals. Besides, polyamides are known
to undergo polymorphic transitions upon heating, and they
may form lamellas with different thickness [37, 46, 47].
0,2
0,1
0,0
–0,1
–0,2
–0,3160 200 240 280
Temperature/ C
Nor
mal
ized
hea
t flo
w/W
g–1
Exo
non-reversing
reversing
PA 4.10PA 6.10PA 10.10PA 10.12
°
Fig. 4 Reversing and non-reversing TOPEM� signals, recordedduring heating for all the bio-polyamides under investigation (at
0.5 �C min-1, amplitude of 1 K, the switching time range of 15–30 s)
Table 3 WAXD peak positions and lattice spacings for bio-polyamides under investigation
Bio-PA Position 2h/� Spacings/nm
001 002 100 010/110 001 002 100 010/110
PA 4.10 8.2 b 11.2 b 20.1 23.7 1.078 0.789 0.438 0.376
PA 6.10 – 9.7 b 20.9 23.8 vw – 0.912 0.425 –
PA 10.10 8.3 b 11.9 b 19.9 23.7 1.058 0.742 0.445 0.375
PA 10.12 8.2 b 11.8 b 20.1 23.1 1.072 0.751 0.441 0.384
Other reflections, position 2h/�
Undefined Smectic b/c Amorphous Undefined
PA 4.10 5.9 20.6 b – 38.9 37.7 40.8
PA 6.10 – 20.4 b 22.4 39.8 37.4 40.5 b
PA 10.10 – 19.7 b 21.7 b 38.6 38.6 40.5
PA 10.12 – 18.3 b 21.4 40.3 37.5 40.5
b Broad reflection, vw very weak
Renewable-sourced bio-polyamides 1231
123
-
The melting point (Tm) increases with increasing amide/
methylene ratio in the polymer backbone (0.17 for PA 4.10
and 0.1 for PA 10.12). This is consistent with the
increasing linear density of hydrogen bonds and was found
also for other PAs [19, 48, 49].
The melting starts actually at temperature lower than Tmand can be attributed to slow melting of the material
crystallized during the first scan in DSC, which was per-
formed to remove thermal history before temperature-
modulated measurements. Upon heating, the less perfect
crystals transform to a supercooled state in which recrys-
tallization takes place, what was observed on the non-re-
versing signal.
The non-reversing component refers processes that are
not reversible in the time frame of the modulation, e.g.,
cold crystallization, melting of the crystallite nuclei, and
recrystallization. In the studied polyamides, with the
exception of PA 6.10, sharp and strong exotherms occur
upon the onset of melting and before the peak melting
temperature. These are assigned to the recrystallization or
so-called crystal perfection. This is the characteristic event
for less perfect crystals, which melt at temperatures below
the thermodynamic melting point and then crystallize to
another crystal type. This is more notably visible for the
PA 4.10, where the recrystallization occurs between the
melting of the two types of crystallites.
For the materials with higher amide concentration (PA
4.10 and PA 6.10), a shallow cold crystallization peak
occurs in the non-reversing signal, well below the onset of
melting. This is related to material that did not have enough
time to crystallize during the preceding cooling scan due to
slow dynamics. As the temperature increases again,
mobility is regained, and the crystallization continues.
Apparently, the crystallization dynamics become slower as
the amide/methylene ratio increases.
The last phenomenon observed in the non-reversing
signal is the final melting of crystal nuclei toward the end
of the reversing melting peak. Due to its small intensity,
this phenomenon is observed only for the high amide/
methylene ratio materials. As the crystalline structure of
PA 6.10 is rather simple, no recrystallization occurs, and
practically only this nuclei melting peak is present in the
DSC curves.
Moreover, when comparing the total heat flow signal
with the reversing component, it can be seen that the latter
one is much higher than the former, indicating exothermic
heat flow of recrystallization and ‘‘crystal perfection.’’
Concerning that the melting of semicrystalline polymer is
followed by recrystallization and reorganization of
remaining crystallites, the total heat flow is the sum of the
endothermic heat flow of melting, the exothermic heat flow
of recrystallization and reorganization, and the contribution
of the heat capacity [50].
The crystallinity-based WAXD measurements corre-
spond well with the one from DSC, except PA 10.12.
Although the values are comparable, a clear correlation in
these data cannot be given. For semicrystalline polymers,
the processing and the thermal history define the thermo-
mechanical features of the material, which in turn deter-
mine, e.g., the relaxation, nucleation, and the crystal
structure and morphology. Therefore, the processed bio-
polyamides differ in their crystal structure from the raw
pellets tested in DSC. A more detailed kinetic study of the
crystallization dynamics with TOPEM� and isothermal
crystallization studies in future work is expected to shed
more light in this issue.
Molecular dynamics
Dynamic mechanical analysis (DMA)
Modulus (storage E0 and loss E00) and tand values of bio-polyamides were recorded as a function of the temperature
for different frequencies of mechanical load. Figure 5
shows a comparative plot of E0 and tand at 5 Hz as afunction of temperature.
The DMA curves show three distinct regions, reflected
as peaks on the tand curves and as downward steps in theE’ curves (Fig. 5). Starting from the low-temperature side,
and following the tand curves, the peak around -150 �Creflects the c-relaxation, which has been attributed tomotions of methylene groups [51]. The composition of the
Table 4 Thermal properties of bio-polyamides by TOPEM� DSC and degree of crystallinity as measured by WAXD
Bio-PA Tcp/ �C Tm/�C DHf/J g-1 vc/%
1 2 Rev Non-rev vc,DSC vc,WAXD
PA 4.10 248 246 251 108.63 49.87 23.11 20.54
PA 6.10 205 221 – 113.33 39.47 29.05 36.97
PA 10.10 199 195 – 223.07 155.34 27.76 30.21
PA 10.12 181 190 – 125.72 75.01 20.78 42.87
Tcp—Temperature of crystallization during ‘‘crystal perfection,’’ Tm—melting point at reversing heat flow curve, DHf—enthalpy of fusionevaluated by integrating the heat flow, vc,DSC—degree of crystallinity in DSC, vc,WAXD—degree of crystallinity in WAXD
1232 J. Pagacz et al.
123
-
bio-polyamides does not affect the timescale of this
relaxation, confirming its local character. We would like to
point out here that a relaxation with very similar timescale
characteristics occurs also in polyethers with methylene
sequences and their polyurethane copolymers [52], thus
supporting the assignment of this relaxation to methylene
sequences, in both systems. The strength of the relaxation,
as quantified by the height of the peaks (Fig. 5), also
supports this assumption.
The next peak around -80 �C is the b-relaxation, typ-ically attributed to non-hydrogen-bonded amide groups
[51] and more specifically to water bound on carbonyl
groups. Within the experimental error, this relaxation
seems unaffected by the length of the methylene sequences
between the amide groups, as a result of its strongly
localized character. The amide bond is very similar to
urethane and urea bonds, having in common the CO–NH
sequence. Indeed, a b-relaxation with similar timescalecharacteristics has been observed in polyurethanes [53], in
which it is typically attributed to the same type of molec-
ular motion (i.e., water molecules bound on carbonyl
groups) [54]. Interestingly, albeit it is observed by dielec-
tric techniques, a similar relaxation occur also in NH2 [55]-
and OH [52]-terminated polyethers, which denotes that the
NH component, too, might have a contribution to this
relaxation.
The prominent peak above 50 �C is typically attributedto the dynamic glass transition [56, 57], although some
authors dispute it [51]. We will come back to this point, in
light of the results by the stochastic modulated DSC results.
Nevertheless, we will refer to the peak as the dynamic glass
transition (a-relaxation). Its peak temperature Ta is con-sidered as a good measure of the calorimetric glass tran-
sition temperature, Tg. In general, the Tg of polyamides
drops with increasing mean amide spacing [58, 59]. Here,
according to DMA, the PA 4.10 variant seems to have
significantly higher Tg than its longer segment counter-
parts, but otherwise, the other bio-PAs do not follow the
trend: With increasing mean amide spacing, the peaks
move to slightly higher frequencies as demonstrated in the
Arrhenius plot of Fig. 6 and in Table 5.
The dynamics of the a-relaxation is known to follow aVogel Fulcher Tamman law [60]:
f ¼ f0 e�B
T�T0 ð3Þ
where f0, B, and T0 are parameters to be determined.
Nevertheless, the frequency range of the DMA is rather
limited for the determination of all three of these param-
eters, so we may assume a pseudo-Arrhenius behavior:
f ¼ f0 e�EactkT ð4Þ
and calculate a pseudo-activation energy upon a linear
fitting on the Arrhenius map. The results are again non-
monotonous with the amide/methylene ratio: The PA 4.10
104
103
–150 –100 –50 0 50 100 150
–150 –100 –50 0 50 100 150
0,15
0,10
0,05
Temperature/ C
Mod
ulus
E ′ /
MP
ata
nδ
f = 5 Hz
γ
γ
β
α
β
α
PA 4.10
PA 6.10
PA 10.10
PA 10.12
°
Fig. 5 Storage modulus E0 and tand curves record with all bio-PAsunder investigation, at 5 Hz. For clarity, only every tenth experi-
mental point has been plotted
1,0
0,5
0,0
–0,5
–1,0
–1,5
–2,0
–2,5
–3,0
2,85 2,90 2,95 3,00 3,05 3,10 3,15 3,20
1000 K / T
Iog(
freq
uenc
y / H
z)
80 75 70 65 60 55 50 45 40
DMA
TOPEM
PA 4.10PA 6.10PA 10.10PA 10.12
Temperature/ C°
Fig. 6 Arrhenius map in the segmental dynamics region of allmaterials under investigation. DMA points correspond to the a-peakat tand, and stochastic modulated DSC (TOPEM�) points correspondto the Richardson midpoint
Renewable-sourced bio-polyamides 1233
123
-
has the lowest value, while for the other bio-PAs the Eact is
a decreasing function of the mean amide spacing, i.e., the
average length of the aliphatic components. However,
always a higher Ta corresponds to lower activation energy.
The discrepancy between the PA 4.10 and other bio-PAs
may be related to different technologies for the production
of these polymers on industrial scale.
Another interesting point is that the traces of all mate-
rials on the Arrhenius map tend to converge toward the
low-frequency/low-temperature region. We will come back
to this point later upon discussion of the stochastic mod-
ulated DSC data in the next section.
The PA 4.10 exhibits higher mechanical modulus, in
both the glassy and rubbery regions, than its other homo-
logs, but this may be related, again, to a different manu-
facturing process. For the PA 6.10, PA 10.10, and PA
10.12, the rubbery modulus is an increasing function of the
average distance between amide bonds. However, the
glassy modulus remains practically unaffected.
Stochastic modulated DSC (TOPEM�)
In Fig. 7, we plot the magnitude c�p (T) of the complex heat
capacity as calculated by the Fourier transformation of the
response to the stochastic modulated stimulus. The form of
this response is similar, from the phenomenological point
of view, to heat capacity steps related to glass transition in
conventional DSC experiments. However, in stochastic
modulated experiments, the temperature location of the
step is a function of frequency: The heat capacity is able to
follow faster perturbations only at higher temperatures,
much similar to what is observed in other types of dynamic
experiments such as DMA or dielectric relaxation spec-
troscopy (DRS).
From the comparative DSC curves (Fig. 7), we cannot
draw any straightforward conclusions regarding the
dependence of glass transition temperature on the structure
of the polyamides; however, we may notice that the heat
capacity step Dcp is an increasing function of the distancebetween amide groups within the PA 6.10, PA 10.10, and
PA 10.12 materials. This may be related to the decreasing
degree of crystallinity, as observed with calorimetry,
leaving more amorphous material to participate in the
segmental dynamics.
As a measure of the glass transition temperature (Tg), we
chose the Richardson midpoint [61]. The Tg values
obtained by this method for each studied frequency are
included in the Arrhenius map in Fig. 6, and as a repre-
sentative value, we have included the Tg (1 mHz) in
Table 5.
Bearing in mind the inaccuracy in determination of the
Tg, the traces are located in the same region, with the
exception of PA 10.10 which has a significantly lower Tg.
Interestingly, the region where the traces are located cor-
responds well with the one defined by the extrapolation of
the DMA traces, bearing in mind the expected concave
functional form of the trace of a cooperative relaxation like
the dynamic glass transition (a-relaxation). This supportsour assignment of the high-temperature peak in the DMA
curves to the dynamic glass transition. Experiments with a
more broadband technique such as dielectric relaxation
spectroscopy are expected to shed more light in the inter-
mediate region between the DMA and TOPEM regions.
The reason why the TOPEM trace of the PA 10.10 is
located in lower temperatures than expected is yet to be
clarified.
In Table 5, we show also the heat capacity step at Tg,
Dcp, as calculated at 1 mHz. The polyamide with theshortest distances between amide groups (PA 4.10) has
rather higher Dcp than its counterparts, which show similarvalues within the experimental inaccuracy.
10 20 30 40 50 60 70 80
Temperature/ C
Hea
t cap
acity
cP*,
tra
nsla
ted
PA 4.10
PA 6.10
PA 10.10
PA 10.12
1 Jg–1K–1
f = 1 mHz
°
Fig. 7 Stochastic modulated DSC curves in the region of glasstransition at 1 mHz. The curves have been translated for clarity
Table 5 Glass transition characteristics and a-relaxation pseudo-ac-tivation energy in the frequency region of DMA
Bio-PA TOPEM DMA
Tg/�C Dcp/J g-1 K-1 Ta/�C Eact/kJ mol-1
PA 4.10 46 0.20 75 270
PA 6.10 45 0.13 61 683
PA 10.10 42 0.14 65 547
PA 10.12 45 0.14 63 532
Tg and Dcp—heat capacity step at glass transition temperatureaccording to Richardson (1 mHz), Ta—a-relaxation tand peak tem-perature (5 Hz), Eact—a-relaxation activation energy in the frequencyregion studied by DMA, calculated from the Arrhenius plot in Fig. 6
1234 J. Pagacz et al.
123
-
Conclusions
Bio-polyamides (bio-PAs) are engineering polymers based
on renewable raw materials that link green chemistry
approach with good mechanical and thermal properties. In
this work, we have performed for the first time the struc-
tural characterization of these new polymers, investigated
their properties in the glass transition region and studied
their molecular dynamics. FTIR analysis revealed that
there are characteristic bands for aliphatic polyamides
which have been assigned to the functional groups vibra-
tions, such as NH, CO, CONH, and CH2. Specific inter-
actions by hydrogen bonding were evidenced by the
analysis of the N–H stretching mode at 3300 cm-1, which
broad profile may suggest that more types of hydrogen
bonds with different bond lengths were formed.
In the X-ray analysis, the characteristic reflection (100)
located at 2h = 20� and corresponding to the interchaindistance in the crystalline structure of the a-crystal formwas observed as predominant phase for all bio-PAs.
Interestingly, for PA 10.12, the intensities of the 100 and
010/110 reflections are approximately the same, and the
weak band can be seen at ca. 21�—this observation mayindicate the coexistence of different crystalline forms, a-and c-triclinic and b-pseudohexagonal one.
The melting point (Tm) was found to increase with
increasing amide/methylene ratio in the polymer backbone,
which is consistent with the decreasing linear density of
hydrogen bonds. The multiple peak of melting in bio-
polyamides can be assigned to the fusion of imperfect
crystals with lower thermal stability, secondary crystal-
lization upon heating and remelting of more perfect
crystals.
Molecular dynamics investigations by dynamic
mechanical analysis (DMA) show three distinct regions
associated with c-relaxation, b-relaxation, and dynamicglass transition. Noteworthy, the composition of the bio-
polyamides does not affect the timescale of the c-relax-ation, confirming its local character. Arrhenius map results
evidence that the PA 4.10 has the lowest pseudo-activation
energy in the 1- to 10-Hz region, while for the other bio-
PAs the Eact is a decreasing function of the mean amide
spacing, i.e., the average length of the aliphatic compo-
nents. However, always a higher Ta corresponds to lower
activation energy. Stochastic modulated DSC results show
that at low frequency/long timescales, the materials with
significantly different amide/methylene ratios have similar
segmental dynamics.
From the methodological point of view, stochastic mod-
ulated DSC provides an insight into the segmental dynamics
at very low frequencies, which agree very well with those of
dynamic mechanical analysis recorded at higher
frequencies. However, this point should be followed in
future work in more detail, also in comparison and cross-
evaluation with more broadband techniques such as dielec-
tric relaxation spectroscopy.
Acknowledgements Authors are grateful to the Polish NationalScience Center for financial support under Contract No. DEC-2011/
01/M/ST8/06834. Authors would also like to thank Dr Daniel Fra-
giadakis, Naval Research Laboratory, USA, who develops and
maintains the software Grafity (distributed free of charge at grafity-
labs.com) which we used for our data analysis.
Open Access This article is distributed under the terms of theCreative Commons Attribution 4.0 International License (http://creati
vecommons.org/licenses/by/4.0/), which permits unrestricted use,
distribution, and reproduction in any medium, provided you give
appropriate credit to the original author(s) and the source, provide a link
to the Creative Commons license, and indicate if changes were made.
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Bio-polyamides based on renewable raw materialsGlass transition and crystallinity studiesAbstractIntroductionExperimentalMaterialsMethodsProcessingStructural characterizationFourier transform infrared spectroscopy (FTIR)Wide-angle X-ray diffraction (WAXD)
Thermomechanical characterizationModulated-temperature DSC (TOPEMreg)DMA analysis
Results and discussionStructural characterization (FTIR)Crystalline structure (WAXD)Melting and crystallization (TOPEMreg DSC)Molecular dynamicsDynamic mechanical analysis (DMA)Stochastic modulated DSC (TOPEMreg)
ConclusionsAcknowledgementsReferences